Method for producing a part made from a superalloy based on nickel and corresponding part

ABSTRACT

A method for manufacturing a blank part in Ni-base superalloy, wherein an alloy is prepared and heat treatments are conducted characterized in that: the said superalloy contains at least a total of 2.5% of Nb and Ta; heat treatment is conducted comprising a plurality of steps: a first step at between 850 and 1000° C. held for at least 20 minutes to precipitate the δ phase at the grain boundaries; a second step held at a temperature higher than the temperature of the first step allowing partial dissolution of the δ phase obtained at the first step; ageing treatment comprising a third step and optionally one or more additional steps at a temperature below the temperature of the first step and allowing precipitation of the hardening phases γ′ and γ″. Part thus obtained.

The invention concerns nickel-base superalloys, and more particularly aheat treatment method which can beneficially be applied to some thereoffor the improvement in particular of their creep resistance and tensilestrength.

By <<nickel-base superalloys>>, is meant alloys in which Ni accounts forat least 50% by weight of the composition (all the percentages given inthis text are weight percentages).

More precisely, the invention concerns a heat treatment methodapplicable to alloys having a content of niobium and tantalum totallingmore than 2.5%, and which are therefore able to cause the onset ofdouble precipitation:

-   -   intergranular precipitation of the δ phase (Ni₃Nb-δ or Ni₃Ta-δ)        at between 800 and 1050° C.;    -   intragranular precipitation of the hardening phases of types        γ′(Ni₃ (Al—Ti)-γ′) and/or γ″ (Ni₃Nb-γ″ or Ni₃Ta γ″) during        ageing conducted at between about 600 and 800° C.

This is particularly the case with the alloy NC19FeNb, commerciallyknown as INCONEL 718® (718) and alloys derived therefrom or comparabletherewith such as 625, 718Plus and 725.

In the aeronautic and land-based gas turbine industry, in which anickel-base resistant alloy has numerous applications, experience hasshown that the fatigue strength of alloys is among the most criticalfactors for the sizing of turbine discs and shafts.

The relatively low cost of the 718 alloy, through the absence of cobaltin its composition and the acquired know-how for its production andtransformation, give this alloy privileged ranking among the highcharacteristic alloys used up to a temperature close to 650° C. However,the increased yield and performance of turbo-machines translates as anincrease in temperature at the output of the combustion chamber, andtherefore requires improved creep resistance of the 718 alloy toincrease the possibilities of extending periods of use up to 650° C. Theimprovement in the creep resistance of the 718 alloy, whilst maintaininga fine grain microstructure (>7 ASTM) so as not to compromise fatiguestrength, is therefore of major industrial interest. It is recalled thatthe ASTM standards governing the estimation of grain size define thegrains as being finer the higher the attributed ASTM number.

Two different thermo-mechanical treatment methods are known andcurrently used to improve the fatigue properties of the 718 alloy.

According to a first option such as described in FR-A-2 089 069, it waschosen to perform thermo-mechanical treatment allowing the Ni₃Nb-δ phaseto be precipitated at the grain boundaries, followed byrecrystallization treatment of the alloy at a temperature below thedissolution temperature of the Ni₃Nb-δ phase, the Ni₃Nb-δ phaseprecipitated at the grain boundaries being used during recrystallizationto prevent grain growth. With this method it is possible to obtainrecrystallized structures with very fine grain size, of ASTM 10 orhigher. Their fatigue characteristics are improved but the creepresistance thereof is insufficient. Indeed it is known that the presenceof the Ni₃Nb-δ phase, having an orthorhombic structure, is detrimentalsince it fixes the niobium and thereby limits the formation of theNi₃Nb-γ″ hardening phase, that is metastable and of centred quadraticstructure. The Ni₃Nb-γ″ hardening phase allows the slowing ofdislocation movement within the crystallographic lattice and therebyimproves creep resistance.

Similarly, it is also known that the presence of the Ni₃Ta-δ phase isdetrimental, since its fixes the tantalum and therefore limits theformation of the Ni₃Ta-γ″ hardening phase.

Another known solution for improving the properties of 718 consists ofconducting ageing directly after thermo-mechanical treatment i.e.without the usual solution heat treatment at between 900 and 980° C.carried out between the thermo-mechanical treatment and the ageingtreatment. Although this option allows limited formation of the Ni₃Nb-δphase which may precipitate during the solution heat treatment, and theobtaining of fine grain size together with improved tensile and fatigueproperties, it does have disadvantages.

It has been found that heterogeneous microstructures are obtained withinone same part owing to major local variations in grain size and to theproportion of δ phase formed during thermo-mechanical treatments.

As a result, creep resistance is degraded compared with prior practiceover a wide temperature and stress range.

Document EP-A-1 398 393 describes treatments of Ni-base superalloys inthe form of directionally solidified single crystals or alloys. If thealloy is a single crystal there is evidently no precipitation of δ phaseat the grain boundaries since there are no grain boundaries. Withdirectional solidification, any precipitation of δ phase could onlyoccur heterogeneously and would not prevent grain growth. At the end ofthe treatment, the grains would be too large in size. In addition, thealloy compositions preferably described in this document would not allowprecipitation of the δ phase, having regard to the Ti, Ta, Nb and Alcontents thereof, since this phase would not be stable on account of thehigh Al content.

Document U.S. Pat. No. 4,459,160 also describes single crystal Ni-basesuperalloys in which no precipitation of δ phase can be observed at thegrain boundaries.

It is the objective of the invention to improve the creep resistance andtensile strength of nickel-base superalloys having a content of niobiumand/or tantalum higher than 2.5%, without deteriorating the fatigueproperties and whilst avoiding the disadvantages of the aforementionedprior art.

For this purpose, the subject of the invention is a method formanufacturing a Ni-base superalloy blank containing at least 50% Ni asweight percentage, according to which an alloy of said superalloy isproduced and the said alloy is subjected to heat treatments,characterized in that:

-   -   the said superalloy, in weight percentage, contains Nb and Ta at        least to a total amount of 2.5%;    -   the said alloy is subjected to heat treatment comprising a        plurality of hold steps distributed as follows:        -   a first hold step during which said alloy is held at between            850 and 1000° C. for at least 20 minutes to precipitate the            δ phase at the grain boundaries;        -   a second hold step during which said alloy is held at a            temperature higher than the temperature of the first hold            step and allowing partial dissolution of the δ phase            obtained at the first step;        -   aging treatment comprising a third step and optionally one            or more additional steps conducted at a temperature lower            than that of the first step and allowing precipitation of            the hardening phases γ′ and/or γ″

Preferably, the Al content of the alloy is equal to or less than 3%.

Preferably, the (Nb+Ta+Ti)/Al ratio of the alloy is 3 or higher.

Preferably, the grain size obtained at the end of the aging treatment ofthe alloy ranges from 7 to 13 ASTM, more preferably from 8 to 12 ASTM,further preferably from 9 to 11 ASTM.

Preferably, the distribution of the δ phase is homogeneous at the grainboundaries on completion of the ageing treatment.

After the second hold step, preferably a quantity of δ phase is obtainedof between 2 and 4%, and is best between 2.5 and 3.5%.

The first and second hold steps are preferably conducted withoutintermediate cooling.

The changeover from the first to the second hold step can then takeplace at a rate of 4° C./min or less, preferably between 1 and 3°C./min.

The first hold step can be conducted at between 900 and 1000° C. for atleast 30 min, and the second hold step at between 940 and 1020° C. for 5to 90 min, the temperature difference between the two temperature holdsbeing at least 20° C.

The content by weight of the alloy may be as follows:

between 50 and 55% nickel,

between 17 and 21% chromium,

less than 0.08% carbon,

less than 0.35% manganese,

less than 1% cobalt,

less than 0.35% silicon,

between 2.8 and 3.3% molybdenum,

at least one of the elements niobium or tantalum, such that the sum ofniobium and tantalum totals between 4.75% and 5.5% with Ta being lessthan 0.2%,

between 0.65 and 1.15% titanium,

between 0.20 and 0.80% aluminium,

less than 0.006% boron,

less than 0.015% phosphorus,

the residual percentage being iron and impurities resulting fromprocessing.

The first hold step can then be conducted at between 920 and 990° C. forat least 30 min, and the second hold step at a temperature of between960 and 1010° C. for 5 to 45 min.

The total content of Nb and Ta of the alloy may then be between 5.2 and5.5%, the first hold step being conducted at between 960 and 990° C. for45 min to 2 h and the second hold step at between 990 and 1010° C. for 5to 45 min.

If the total content of Nb and Ta of the alloy is between 4.8 and 5.2%,the first hold step can be conducted at between 920 and 960° C. for 45min to 2 h and the second hold step can be conducted at between 960 and990° C. for 5 to 45 min.

The content by weight of the alloy may be:

between 55 and 61% nickel,

between 19 and 22.5% chromium,

between 7 and 9.5% molybdenum,

at least one of the elements niobium or tantalum, such that the sum ofniobium and tantalum totals between 2.75 and 4% with Ta being less than0.2%,

between 1 and 1.7% titanium,

less than 0.55% aluminium,

less than 0.5% cobalt,

less than 0.03% carbon,

less than 0.35% manganese,

less than 0.2% silicon,

less than 0.006% boron,

less than 0.015% phosphorus,

less than 0.01% sulphur,

the residual percentage being iron and impurities resulting fromprocessing.

The alloy may have a weight content of:

between 12 and 20% chromium,

between 2 and 4% molybdenum,

at least one of the elements niobium or tantalum, such that the sum ofniobium or tantalum is between 5 and 7% with Ta less than 0.2%,

between 1 and 2% tungsten,

between 5 and 10% cobalt,

between 0.4 and 1.4% titanium,

between 0.6 and 2.6% aluminium,

between 6 and 14% iron,

less than 0.1% carbon,

less than 0.015% boron,

less than 0.03% phosphorus

the residual percentage being nickel and impurities resulting fromprocessing.

Preferably the aforementioned alloys, in weight percentage, have aphosphorus content of more than 0.007%.

In general, the first and the second hold steps can be conducted at δphase sub-solvus temperatures of the alloy, the first hold step beingconducted at a temperature of between 50° C. below the δ solvustemperature and 20° C. below the δ solvus temperature, and the secondhold step being conducted at a temperature of between 20° C. below the δsolvus temperature and the δ solvus temperature.

The temperature of the hot-formed blank can be held constant during atleast one of the steps.

The said third step can be conducted at between 700 and 750° C. for 4 to16 h and a fourth step is then conducted at between 600 and 650° C. forbetween 4 and 16 h, cooling at 50° C./h to +/−10° C./h being conductedbetween said third and fourth steps.

Between the first and second steps, it is possible to hold thehot-formed alloy at an intermediate temperature between the temperaturesof the first and second steps for a maximum time of 1 h.

The said blank may have been produced in the form of an ingot and thenhot-worked.

The said blank may have been produced using a powder metallurgy method.

A further subject of the invention is a part in a nickel-basesuperalloy, characterized in that it was obtained from a blank producedaccording to the above-mentioned method.

This may be a blank of a part for an aeronautic or land-based gasturbine.

As will have been understood, the invention consists of subjecting aNi-base alloy containing Nb and/or Ta to a heat treatment for whichstructural hardening is obtained by precipitation of the gamma′(Ni₃Ti-γ′) and/or gamma″ (Ni₃Nb-γ″ and/or Ni₃Ta-γ″) hardening phases,these phases respectively comprising Titanium and Niobium and/orTantalum. The heat treatment comprises at least three hold steps whichchronologically are the following:

-   -   a first hold step conducted at 850-1000° C. which is intended to        precipitate the delta phase Ni₃Nb-δ and/or Ni₃Ta-δ at the grain        boundaries, with substantially homogeneous distribution of this        phase in the grain boundaries, and to homogenize the        microstructure of the material; with regard to partly        recrystallized microstructures it also allows completion of        recrystallization and causes the δ phase to precipitate at the        boundaries of the new recrystallized grains;    -   a second hold step conducted at a temperature higher than that        of the first step and intended for part dissolution of the said        delta phase Ni₃Nb-δ and/or Ni₃Ta-δ, whilst maintaining the        substantially homogeneous distribution obtained after the first        step, and avoiding grain enlargement; the second step is        completed by oil quench or air cooling;    -   the third step and any optional following steps are thermal        ageing steps conducted at a temperature below the temperature of        the first step and allowing precipitation of the gamma′        (Ni₃(Al—Ti)-γ′) and/or gamma″ (Ni₃Nb-γ″ or Ni₃Ta-γ″) hardening        phases.

One or more intermediate cooling operations are possible between eachstep but are not compulsory.

The method of the invention allows parts to be produced which, comparedwith those of the prior art having the same composition, offer a bettercompromise between yield strength under heavy loading, high fatiguestrength and long creep resistance lifetime.

The invention will be better understood on reading the followingdescription given with reference to the following appended figures:

FIGS. 1 to 3 which schematize three examples of the two first heattreatment steps according to the invention, FIG. 2 also showing anintermediate step between the first and second steps; the temperaturesalong the Y-axis are referenced in relation to δ phase solvustemperature.

FIGS. 4 to 9 which give micrographs of alloys subjected to referenceheat treatments (FIGS. 4 to 7) and according to the invention (FIGS. 8,9).

The method for manufacturing a part in Ni superalloy according to theinvention may be initiated by preparing and casting an ingot of saidsuperalloy using conventional methods such as a double melt method (VIMVacuum Induction Melting—VAR Vacuum Arc Remelting) or triple melt(VIM—ESR (Electroslag remelting)—VAR). However, the method of theinvention can also be applied to a blank produced by power metallurgy.In the remainder of the text the examples of applications described areexamples in which the starting product is obtained by the conventionalroute called <<ingot metallurgy>>, but the transposition thereof topowder metallurgy will be obvious for persons skilled in the art. Thetreatments following after hot-working that are characteristic of theinvention will be the same in both cases.

The initial microstructure of a product (on the understanding that theterm <<product>> designates a semi-product or blank of a part) beforethe treatment typical of the invention, may vary in relation to thedeformation thermo-mechanical treatments conducted upstream, for exampleforging, punching or hot rolling:

-   -   metallurgical state 1 (or <<state 1>>): the delta phase Ni₃Nb-δ        and/or Ni₃Ta-δ can be present at the grain boundaries but not        uniformly distributed between the grains subsequent to        deformation conducted at a temperature lower than δ phase solvus        temperature;    -   metallurgical state 2 (or <<state 2>>): the delta phase Ni₃Nb-δ        and/or Ni₃Ta-δ may be absent or practically absent (<1%) from        the microstructure subsequent to deformation conducted for        example at a temperature higher than the δ phase solvus.

In the first case, i.e. starting from metallurgical state 1, the firsttreatment step according to the invention allows the distribution of theδ phase to be homogenized within the microstructure, and permits thereducing of local variations in the δ phase fraction present after thethermo-mechanical treatments due to temperature variations of greater orlesser extent after deformation. Persons skilled in the art, ifnecessary, are easily able to adjust the parameters for conducting thefirst step through routine testing, in order to optimize thishomogenization of δ phase distribution.

In the second case i.e. starting from metallurgical state 2, the firsttreatment step according to the invention allows (substantially)homogeneous precipitation of δ phase at the grain boundaries which weredevoid of this phase after the thermo-mechanical treatment. Personsskilled in the art may also, through routine testing, adjust theparameters if necessary for conducting the first step so as to optimizethis homogenization of δ phase distribution.

Whether in the first or second case, the first step also allowscompletion of recrystallization in the regions where recrystallizationmay not have been complete during thermo-mechanical treatment, and itthereby homogenizes the global structure of the alloy.

At the second step of the treatment according to the invention,conducted at a temperature close to the δ phase solvus, the delta phaseNi₃Nb-δ and/or Ni₃Ta-δ is partly dissolved.

At the second step, the dissolution of the δ phase takes place insubstantially uniform manner. The so-called residual δ phase i.e. thenon-dissolved δ phase maintains the same distribution as obtained afterthe first step. On this account, the residual δ phase remainssubstantially uniformly distributed around the grains, allowing theslowed growth of all the grains and the limiting and even the avoidingof the onset of large grains at the second step, which is conducted at atemperature higher than that of the first step. The homogeneousdistribution of the δ phase at the grain boundaries promotes thehomogeneity of grain size in the microstructure of the alloy at the endof treatment.

The second step therefore allows a reduction in the quantity of δ phaseobtained after the first step, down to a residual quantity that isoptimally lower than 4%, even below 3.5% whilst avoiding grainenlargement.

The greater dissolution of the δ phase in a fine-grain, homogeneousmicrostructure allows more niobium to be released for precipitation ofthe gamma′ and/or gamma″ hardening phases during a third step and evenother subsequent steps forming an ageing treatment of the alloy.

In unexpected manner, the inventors have found that the absence of thefirst treatment step does not allow these effects to be obtained,irrespective of the initial microstructure after thermo-mechanicaltreatment.

Evidently, for an initial microstructure devoid of δ phase (state 2) theabsence of the first step will not allow homogenization of the globalstructure of the material and precipitation of the δ phase at the grainboundaries, limiting subsequent growth of the grains during the secondstep.

In the absence of the first step, when the initial microstructureresults from sub-solvus deformation which led to δ phase precipitation(state 1), the distribution of the δ phase is heterogeneous (see FIGS. 4and 5). Therefore some grains may contain a large quantity of δ phase atthe grain boundaries, or little or no δ phase at the grain boundaries,or even a heterogeneous distribution of δ phase at the grain boundaries.

By conducting a heat treatment directly at the temperature of the secondstep, without a temperature hold at the temperature of the first step,the grains which are not surrounded with δ phase or which have little δphase at the grain boundaries, or a non-uniformly distributed δ phasewill enlarge uncontrollably up to a grain size possibly exceeding aboutASTM 5-6. The presence, even much localised presence of ASTM 5-6 grains(see FIGS. 6 and 7), reduces the fatigue lifetimes by a factor of 10compared with a homogeneous microstructure having ASTM size 10 grains.The combination of the first and second steps according to the inventiontherefore (see FIGS. 8 and 9) allow the partial dissolution of the δphase and in homogeneous manner, whilst avoiding the presence of theselarge ASTM 5-6 grains which is redhibitory for guaranteeing high fatigueproperties.

With an initial microstructure comprising δ phase (state 1), the absenceof the first step does not therefore allow the desired microstructure tobe obtained i.e. having a residual, homogeneous δ phase contentpreferably less than 4% with a homogeneous and acceptable grain size.

The preferred grain size for the products derived from the method of theinvention follows from the desire to achieve a good compromise betweenconflicting properties with regard to their grain size requirements.Fatigue strength and tensile strength effectively benefit from a smallgrain size, whereas creep resistance and crack resistance benefit from acoarse grain size. In this perspective, the preferred grain sizes areASTM 7 to 13, preferably ASTM 8 to 12, and best ASTM 9 to 11.

The absence of the second step after conducting the first stepcorresponds to treatments of usual type performed on superalloy productsto which the invention applies, and for which it was seen above thatthey are not satisfactory.

In addition, for an initial microstructure devoid of δ phase (state 2),and if neither of the two first steps required by the invention areconducted, and therefore if thermal ageing treatment is directly appliedto the alloy (so-called <<Direct Aged>> treatment) after its hot-workingat δ phase super-solvus temperature (state 2), in the final structure atotal absence of δ phase is obtained which is undesirable.

Unexpectedly, the inventors were effectively able to evidence that apresence of δ phase of between 2 and 4% and optimally between 2.5 and3.5% allows the properties of the material to improved without weakeningthereof.

On the other hand, microstructures which are devoid of δ phase are ingeneral more subject to intergranular weakening which considerablyreduces high temperature ductility and strongly increases the alloy'ssensitivity to notch effect (for example premature creep rupture at anotched point). Therefore, when the δ phase is absent afterthermo-mechanical treatment, the first step is also necessary to createa minimum amount of δ phase distributed homogeneously at the grainboundaries and to homogenize the global structure of the material.

The hold period of the alloy at the temperature of the first step isequal to or more than 20 minutes. The temperature of the first step isbetween 850 and 1000° C. to precipitate the δ phase. The temperature andthe holding time are adjusted in relation to the heterogeneity of themicrostructure after deformation, and with a view to maintaining anamount of δ phase after the second step that is higher than the minimumrequired for hot ductility.

The second step conducted at a temperature higher than the first step istherefore necessary to allow lowering of the quantity of δ phase bydissolution down to the desired level, preferably to a content ofbetween 2 and 4% and optimally between 2.5 and 3.5%, to release the Nband/or Ta needed for precipitation of the γ′ phase and/or γ″ phasewhilst maintaining a sufficient quantity of Nb and/or Ta in δ phase formdistributed homogeneously around the grains for the hot ductility of thematerial.

The temperature and the duration of the second step are adjusted inrelation to the fraction of δ phase obtained after the first step inorder to obtain the desired residual fraction of δ phase, whilstavoiding grain enlargement. The duration of the second step is alsorelated to the temperature determined for this step. In general, theduration of the second step is shorter the higher the temperaturethereof.

According to one preferred variant of the invention, the two firsttreatment steps are successive steps (FIGS. 1 and 2).

By <<successive treatment steps>>, is meant that the changeover from thefirst treatment step to the second treatment step takes place byprogressively increasing the temperature to move from the first steponto the second without passing through an intermediate temperaturelower than that of the first step.

The succession of the two first steps without descending down to atemperature lower than that of the first step, for example down toambient temperature, allows the avoiding of large temperature gradientsinside the treated sample, and the avoiding of heterogeneous dissolutionof the δ phase which could cause grain enlargement in some regions. Itis therefore preferable to adopt a sufficiently low rate of temperaturerise (<4° C./min) between the steps so that the temperature remainshomogeneous within the treated sample during the second step. It wasverified at the second step that the temperature was homogeneous after 5minutes within a cylindrical sample of 1000 cm³ after a rate oftemperature rise of 2° C./min from the first step. Therefore, anychangeover between the two steps at a temperature lower than the firststep risks increasing the time needed for homogenization of thetemperature within the sample during the second stage, and riskspromoting heterogeneous dissolution of the δ phase. Nevertheless, saidchangeover to a temperature lower than that of the first step is notexcluded by the invention (FIG. 3) if, in particular in relation to thesize of the treated part, the parameters of the second step areadjusted, optionally by adding an intermediate step so as to avoid thepossible disadvantages that have just been mentioned.

Preferably, the first treatment step is conducted at a temperature ofbetween about 900 and 1000° C. for a time of at least 30 minutes, andthe second treatment step is conducted at a temperature higher than thatof the first step at between 940° and 1020° C. for a time of between 5and 90 minutes. The difference in temperature between the two steps musttherefore be at least 20° C. The temperature ranges and time durationsthus obtained allow a homogeneous microstructure to be obtained with anadequate grain size i.e. between ASTM 7 and 13, preferably between ASTM8 and 12, best between ASTM 9 and 11, and a residual δ phase fraction ofbetween 2% and 4%.

As will have been understood the invention is firstly based on a synergyeffect between the two first steps, and the optimised balancing betweenthese two first two steps allows the objectives set by the invention tobe best met.

The δ phase solvus temperature is directly dependent upon theniobium+tantalum content of the alloy. The quantity of niobium and/ortantalum present in the composition of the alloy therefore has a directinfluence on the temperature and duration of each step.

If an alloy of type 718 is used (whose standardized composition isdetailed below), it is indicated to conduct the first temperature holdat between 920 and 990° C. for at least 30 min, and the secondtemperature hold at between 960 and 1010° C. for 5 to 45 min. Theoptimal durations of the treatments are also dependent upon themassiveness of the part to be treated, and can be determined usingmodelling or experiments usually used by those skilled in the art.

For a total Nb and Ta content of the 718 alloy (with less than 0.2% Ta)of between 5.2 and 5.5%, the first step is preferably conducted at atemperature of between about 960° C. and 990° C. for a time of betweenabout 45 minutes and 2 hours, and the second step is preferablyconducted at a temperature of between about 990° C. and 1010° C. for atime of between about 5 and 45 minutes.

For a Nb+Ta content of the 718 alloy (with less than 0.2% Ta) of betweenabout 4.8 and 5.2%, the first step is preferably conducted at atemperature of between about 920° C. and 960° C. for a time of betweenabout 45 minutes and 2 hours, and the second step is preferablyconducted at a temperature of between 960° C. and 990° C. for a time ofbetween about 5 and 45 minutes. The duration of treatment is alsodependent upon the massiveness of the part to be treated.

The temperatures at the treatment steps are generally held substantiallyconstant throughout the duration of the temperature hold.

The rate of temperature rise between the first and second step ispreferably lower than 4° C./min, to avoid temperature gradients that aretoo high, especially if the parts being treated are large parts.

The rate of temperature rise from the first to the second step ispreferably between 1° C./min and 3° C./min.

The invention therefore applies to nickel-base superalloys containing atleast 50% Ni, in which the sum of Nb+Ta exceeds 2.5% by weight.

In one particular case, the alloy is nickel-base alloy of 718 type alsocalled NC19FeNb (AFNOR standard), with a weight content of,

between 50 and 55% nickel,

between 17 and 21% chromium,

less than 0.08% carbon,

less than 0.35% manganese,

less than 0.35% silicon,

less than 1% cobalt,

between 2.8 and 3.3% molybdenum,

at least one of the elements niobium or tantalum, such that the sum ofniobium and tantalum is between 4.75 and 5.5% with Ta less than 0.2%,

between 0.65 and 1.15% titanium,

between 0.20 and 0.80% aluminium,

less than 0.006% boron,

less than 0.015% phosphorus,

the residual percentage being iron and impurities resulting fromprocessing.

The elements for which no minimum content is given may only be presentin trace form, in other words at a content which may be zero, at allevents sufficiently low to have no metallurgic effect (this is true forthe compositions which will be mentioned).

Advantageously, an addition of phosphorus allows grain boundary strengthto be reinforced, in particular against stresses such as creep andnotched creep. The application of the invention to said alloy withphosphorus content higher than 0.007% and lower than 0.015% is ofparticular interest since the creep gain obtained is distinctly greater.It therefore becomes easily possible to improve creep lifetimes by afactor of 4 whilst maintaining the same grain size. This presence ofphosphorus, for the same reasons, may also be recommended for the otherexamples of alloy give below.

In another particular case, the alloy is a nickel-base superalloy of 725type, having a weight content of:

-   -   between 55 and 61% nickel,    -   between 19 and 22.5% chromium,    -   between 7 and 9.5% molybdenum,    -   at least one of the elements niobium or tantalum, such that the        sum of niobium and tantalum is between 2.75 and 4% with Ta less        than 0.2%, between 1 and 1.7% titanium,    -   less than 0.55% aluminium,    -   less than 0.5% cobalt    -   less than 0.03% carbon,    -   less than 0.35% manganese,    -   less than 0.2% silicon,    -   less than 0.006% boron,    -   less than 0.015% phosphorus,    -   less than 0.01% sulphur,    -   the residual percentage being iron and impurities resulting from        processing.

In another particular case, the alloy is a nickel-base superalloy of718PLUS type, with a weight content of:

-   -   between 12 and 20% chromium,    -   between 2 and 4% molybdenum,    -   at least one of the elements niobium or tantalum, such that the        sum of niobium or tantalum is between 5 and 7%, with Ta lower        than 0.2%,    -   between 1 and 2% tungsten,    -   between 5 and 10% cobalt,    -   between 0.4 and 1.4% titanium,    -   between 0.6 and 2.6% aluminium,    -   between 6 and 14% iron,    -   less than 0.1% carbon,    -   less than 0.015% boron,    -   less than 0.03% phosphorus

the residual percentage being nickel and impurities resulting fromprocessing.

In general, the alloy is a nickel-base superalloy characterized by acontent of niobium+tantalum higher than 2.5% and by the presence of anintergranular phase of Ni₃Nb—Ta type (δ phase) at between 800° C. and1050° C. and by the presence of an intragranular phase ofNi₃(Al—Ti)-(γ′) type and/or of Ni₃Nb—Ta (γ″) type at between 600 and800° C. For a nickel-base superalloy containing more than 2.5% niobiumand/or tantalum and characterized by the presence of an intergranularphase containing niobium and/or tantalum and of Ni₃Nb—Ta type, theeffect of the invention is also found even in the absence of the γ″hardening phase Ni₃Nb—Ta. The greater dissolution of the intergranularphase of delta Ni₃Nb—Ta type therefore releases niobium (γ′-gen element)which inserts itself in solid solution in the γ′ hardening phase—Ni₃(Al, Ti) and hardens the latter.

The treatment of the invention may comprise a fourth step allowingcompletion of the precipitation of the gamma″ (Ni₃Nb—Ta-γ″) and/orgamma′ (Ni₃(Al—Ti)-γ′) hardening phases at a temperature lower than thetemperature of the third step.

For example, provision can be made for a third step at between 700 and750° C. for 4 h to 16 h followed by cooling at 50° C./h+/−10° C./h downto the temperature of the fourth step at between 600° C. and 650° C.which is held for between 4 h and 16 h.

The treatment of the invention may also comprise at least oneintermediate step of short duration (no more than 1 h, see FIG. 2)between the first step and the second step to facilitate homogenisationof the temperature within parts of large-size during the temperaturerise between the two first steps.

According to the invention, in which the (Ta+Nb) content of the alloy isat least 2.5%, it is recommended that the Al content does not exceed 3%,so as not to cause precipitation of the γ′ phase at the grainboundaries. Over and above 3% Al, the γ′ phase tends to stabilise to thedetriment of the δ phase and the Nb comes to insert itself in the γ′phase.

Also, still to give priority to the precipitation of the δ phase at thegrain boundaries, it is preferable that the (Nb+Ta+Ti)/Al ratio shouldbe 3 or higher.

The invention will now be illustrated by several examples of embodimentof the heat treatment according to the invention, given in non-limitingmanner.

The first examples of embodiment of the method according to theinvention are applied to articles in 718 alloy obtained afterthermo-mechanical treatment of an alloy obtained via conventional routeof VIM+VAR+forging, but which could just as well have been obtained bypowder metallurgy, and typically intended for the manufacture ofaeronautic turbine discs.

On experimental level, using a VIM method we prepared then re-meltedingots in 718 alloy using the VAR method which were then hot-workedaccording to three different schedules of thermo-mechanical treatment(TTM, cf Table 2) numbered 1 to 3 in Table 2. The products obtainedafter thermo-mechanical treatment were cut up into samples (designated Ato P in Table 1). The samples were then subjected to different heattreatments (TTH) comprising two to four steps depending on differentcases (see Table 2).

The schedule for thermo-mechanical treatment N° 1 comprised rollingconducted with different passes at a temperature higher than the δ phasesolvus temperature of the alloy. The products formed according tothermo-mechanical treatment schedule N° 1 are bars whose metallurgicalstructure is devoid of delta phase (metallurgical state 2). In Table 2the samples F, K, L, N were produced from bars obtained according tothis first thermo-mechanical treatment schedule.

The thermo-mechanical treatment schedule N° 2 was a conventional forgingschedule with reheat (by <<heat>> is meant holding in a furnace followedby deformation; <<reheat>> therefore means two deformation steps, eachbeing preceded by holding in the furnace) at a temperature lower thanthe δ phase solvus temperature of the alloy (<<sub-solvustemperature>>). This schedule allows precipitation of the δ phase in thealloy. The products formed according to the thermo-mechanical treatmentschedule N° 2 are so-called pancakes (a product globally in the roughshape of a flat disc resulting from deformation by forging) whosemetallurgical structure contains some δ phase distributedheterogeneously at the grain boundaries (metallurgical state 1, seeFIGS. 4 and 5). In Table 2, the samples C, E and H were produced frompancakes obtained according to this second thermo-mechanical treatmentschedule.

The thermo-mechanical treatment range N° 3 was a conventional stampingschedule in a single heat step at a temperature lower than the δ phasesolvus of the alloy. The products formed according to thethermo-mechanical treatment schedule N° 3 are blanks of discs whosemetallurgical structure contains some δ phase distributed in highlyheterogeneous manner at the grain boundaries (Metallurgy state 1, seeFIGS. 4 and 5). In Table 2, the samples A, B, D, G, I, J, M, O and Pwere prepared from blanks of turbine discs obtained according to thisthird thermo-mechanical treatment schedule.

Samples A to P were then subjected to five different schedules of heattreatment (<<TTH>>) designated a, b, c, d, e (TTH column, in Table 2)comprising two to four steps as per each case.

The heat treatment schedules of types <<a>> or <<b>> are reference heattreatment schedules representing the state of the art.

The heat treatment schedules of type <<a>> comprise one step ofso-called isothermal solution treatment and two ageing steps. For theseschedules, the solution heat treatment for samples A, B, C, D, F and Pconsisted of holding the alloy at a constant temperature of between 955and 1010° C. for 40 to 90 minutes. The two ageing steps consisted of onehold at 720° C. for 8 hours followed by controlled cooling at 50° C./hdown to a temperature hold of 620° C. for 8 hours.

The type <<b>> heat treatment schedule known as <<Direct Aged>> does notcomprise any solution heat treatment and consists solely of two ageingsteps conforming to type <<a>> treatments. Only sample E was subjectedto the type <<b>> schedule.

The type <<c>> heat treatment schedules are in conformity with theinvention and comprise two so-called solution heat treatment steps,respectively designated as the 1^(st) step and 2^(nd) step, and one ortwo ageing steps, respectively designated as the 3^(rd) step and 4^(th)step.

For these schedules which concerned samples G, H, J, K, M and N, the1^(st) solution heat treatment step consisted of holding the alloy at aconstant temperature of between 940° C. and 980° C. for about 50 to 60minutes. The 2nd solution treatment step consisted of holding the alloyat a constant temperature of between 980° C. and 1005° C. for about 15to 40 minutes. The changeover from the 1^(st) temperature hold to the2^(nd) temperature hold was performed by controlled reheating at a rateof about 2° C./min. The 3^(rd) and 4^(th) ageing steps were inconformity with the corresponding ageing steps of the type <<a>>reference schedules except for samples H and J.

With regard to sample H, the temperature of the 3^(rd) ageing treatmentstep was brought to 750° C. instead of 720° C. as used for the othersamples. This difference allowed the demonstration that the field of theinvention is not limited to restricted temperature conditions andduration of ageing steps, but on the contrary that the invention canalso be applied with temperatures and durations of ageing steps such asthose used in the field of nickel-base superalloys.

As for sample J, this sample was only subjected to one ageing step at720° C. for 10 hours. The ageing treatment applied to sample J showsthat invention can also be applied when the alloy only undergoes asingle ageing treatment step.

The type <<d>> schedules of heat treatments comprised two solution heattreatment steps and two ageing steps. Samples I and L were treated inaccordance with these schedules. However these treatments do not conformto the invention on account of a second step conducted at a temperaturethat is too high or for duration that is too long. The conditions of the2^(nd) step effectively cause too extensive dissolution of the δ phaseand grain growth is no longer controlled, leading to major, uncontrolledgrain enlargement during the second step for samples I and L.

The type <<e>> heat treatment schedule comprised a single solution heattreatment step at 1005° C. for 15 minutes, and two ageing steps. Onlysample O was obtained according to this heat treatment schedule whichdoes not conform to the invention as explained below.

Samples A to L and O were alloys of type 718 with 5.3% of Nb and 40 ppmof P. Sample N was an alloy of type 718 with 5.0% Nb and 40 ppm of P.Samples M and P were alloys of type 718 with 5.3% Nb and 80 ppm of P.

TABLE 1 compositions of tested samples Samples Ni % Fe % Cr % Al % Ti %Nb % Mo % B % C % P % A-L, O 54.2 resid. 17.9 0.5 0.97 5.3 3 0.003 0.030.004 N 53.7 resid. 17.9 0.49 0.98 5.0 3 0.003 0.02 0.004 M, P 54.0resid. 18.1 0.5 1.00 5.3 3 0.003 0.03 0.008

Table 2 summarizes the treatment conditions for the different samples,and the ASTM grain sizes and percentages of surface δ phase which can beseen in a micrograph.

Table 3 summarizes the main mechanical properties of some of thesesamples, namely:

-   -   the yield strength (YS) during a tensile test at 20° C.    -   the ultimate tensile strength during a tensile test at 20° C.        (UTS)    -   the number of cycles before rupture during a fatigue test at        450° C., comprising, with sinusoidal cycle and maximum stress of        1050 MPa, a frequency of 10 Hz and a load ratio R of 0.05;    -   the lifetime during a creep test at 650° C. under a stress of        550 MPa and under a stress of 690 MPa.

The grain size is defined according to the ASTM standard and if thegrain size is relatively inhomogeneous, the maximum grain size (ALA) isalso specified.

TABLE 2 Characteristics and treatments of the different test samplesSolution heat treatment steps Ageing steps 1^(st) step 2^(nd) step3^(rd) step cooling 4^(th) step Microstructure Time Ramp Time T° Time T°Time Grain Sample Alloy Nb % TTM TTH T° (° C.) (min.) ° C./min T° (° C.)(min.) (° C.) (h.) ° C./h (° C.) (h.) ASTM % δ A 718 5.3 3 a 955 60 — —— 720 8 50 620 8 11-12 5.9 B 718 5.3 3. a 970 60 — — — 720 8 50 620 811-12 5.1 C 718 5.3 2. a 975 90 — — — 720 8 50 620 8 10 4.8 D 718 5.3 3.a 1010  40 — — — 720 8 50 620 8  9 2   ALA 5 E 718 5.3 2. b — — — — —720 8 50 620 8 10-14 3-5.5 F 718 5.3 1 a 970 60 — — — 720 8 50 620 8 9-10 5.5 G 718 5.3 3 c 980 60 2 1005 15 720 8 50 620 8 11-12 3.1 H 7185.3 2 c 980 60 2 1005 15 750 8 50 620 8 10 2.9 I 718 5.3 3 d 970 80 21005 60 720 8 50 620 8  9 1.9 ALA 5 J 718 5.3 3 c 970 50 2 995 40 720 10— — — 11-12 3.5 K 718 5.3 1 c 980 60 2 1000 20 720 8 50 620 8  9-10 3.3L 718 5.3 1 d 975 60 2 1015 20 720 8 50 620 8 8-9 1.4 ALA 4 M 718 5.3 3c 980 60 2 1005 15 720 8 50 620 8 11-12 3.0 N 718 5.0 1 c 940 60 2 98015 720 8 50 620 8  9-10 3.4 O 718 5.3 3 e 1005 15 720 8 50 620 8 10 3.2ALA 5 P 718 5.3 3 a 970 60 — — — 720 8 50 620 8 11-12 5.3

The products in 718 alloy, F, K, L, N were therefore transformedaccording to thermo-mechanical schedule n° 1 which does not allow δphase precipitation.

Product F is a reference sample which, after thermo-mechanical treatmentschedule n° 1, was treated using standard type <<a>> thermo-mechanicaltreatment of alloy 718 (treatment comprising a single solution heattreatment step at δ phase sub-solvus).

Product L was treated with two-step solution heat treatment but with asecond step conducted at too high temperature and with too longduration, outside the field of the invention for a 718 alloy.

Products K and N do not have the same niobium content, but both weresubjected to a heat treatment schedule of type <<c>> according to theinvention.

The products in 718 alloy identified as C, E and H were transformed asper the thermo-mechanical schedule n° 2 which allows heterogeneousprecipitation of the δ phase.

Product C is a reference sample which, after the thermo-mechanicalschedule n° 2, was treated as per standard <<a>> type heat treatment ofalloy 718 (treatment comprising only one solution heat treatment atsub-solvus temperature).

Product E is also a reference sample which, after thermo-mechanicalschedule n° 2, was treated as per type <<b>> heat treatment schedule andwas therefore directly aged after forging and therefore did not undergosolution heat treatment before ageing.

Product H was subjected to heat treatment according to the invention(type <<c>>) with two-step solution heat treatment within the field ofthe invention.

The products in 718 alloy identified as A, B, D, G, I, J, M, O and Pwere transformed according to thermo-mechanical schedule n° 3 whichallows highly heterogeneous precipitation of the δ phase.

After thermo-mechanical treatment n° 3, the products A, B and P weretreated as per the standard treatment for alloy 718 (treatment of type<<a>> comprising a single sub-solvus solution heat treatment).

Product D was treated with treatment comprising only one solution heattreatment step but at higher temperature than for products A, B and P,i.e. at a temperature close to δ phase solvus.

After thermo-mechanical treatment, product I was subjected to two-stepsolution heat treatment but whose duration for the second step was toolong having regard to the temperature. The heat treatment applied toproduct I therefore lies outside the field of the invention.

After thermo-mechanical treatment n° 3, product G was treated withtwo-step solution heat treatment within the field of the invention (heattreatment <<c>>). Product J was also treated with two-step solution heattreatment within the field of the invention, but was not treated with afourth step.

Product M was treated with two-step solution heat treatment within thefield of the invention, but has a phosphorus content of 0.008% which istwice higher than that of products A-L and N-O.

Product O was subjected to <<e>> heat treatment with one-step solutionheat treatment; this treatment lies outside the field of the invention.

Product P is a reference sample having a phosphorus content of 0.008%.It was treated using a standard treatment schedule for alloy 718(treatment of <<a>> type comprising a only one solution heat treatmentat sub-solvus temperature).

Products A, B, C which were treated with standard sub-solvus heattreatment (schedule type <<a>>) have a fine grain microstructure (>9ASTM) but comprise a fraction of δ phase (>4.5%) greater than thepreferably desired δ phase fraction according to the invention. Themechanical properties obtained with these products constitute thereference for assessment of the tensile, fatigue and creep propertiesobtained with the thermo-mechanical schedules (TTM) 2 and 3.

TABLE 3 Mechanical properties of the tested samples Fatigue 450° C. R =0.05 Tensile 20° C. 10 Hz σmax = Creep lifetime (h) UTS 1050 MPa 650° C.Sample YS (MPa) (MPa) Lifetime (cycles) 550 MPa 690 MPa A 1210 1470 >3000 000 290 40 B 1240 1480 >3 500 000 340 60 C 1170 1465   1 400 000 35070 D 1310 1495      60 000 1600 200 E 1350 1520 >3 000 000 180 40 G 13401520 >3 000 000 940 120 H 1290 1505 >3 000 000 770 150 M 1335 1520 >3000 000 1400 330 P 1245 1492 >3 000 000 500 80

Product D was treated at a higher temperature than products A, B, and C,it comprises ASTM 5 size grains and a δ phase distributedheterogeneously (<2.5%) that is lower than the preferably desired δphase according to the invention. It is ascertained that this treatmentdid not allow a fine grain microstructure to be maintained (at leastASTM 7, preferably at least 8, and best ASTM 9) nor the satisfactoryfatigue properties found for products A, B, and C. The considerablereduction in the fatigue lifetimes is attributable to the presence oflarge-size ASTM 5 grains which form fatigue initiator sites.

Product E which was directly aged after thermo-mechanical treatment N° 2has a very heterogeneous grain size (ASTM 10 to 14) and major variationsin the level of δ phase, this level being found in most of the regionsof the part (in particular the regions subject to creep) and is higherthan the desired fraction of δ phase. Although the tensile and fatigueproperties of product E are greater than those of products A, B, C, itis ascertained that the creep lifetimes obtained with product E areshorter than the creep lifetimes for products A, B, C.

The absence of solution heat treatment does not allow homogenization ofthe microstructure and is responsible for the presence of very finegrains (>12 ASTM) and of δ phase fractions that are too large, which arethe cause of this degradation of the creep properties.

The absence of solution heat treatment for product E also allows themaintaining of residual work hardening after forging, which isbeneficial for tensile properties but is detrimental to creep resistancein the low stress area.

Products G, H, M were treated within the field of the invention and havea fine grain microstructure (>9 ASTM) and a δ phase fraction (2.9% and3.5%) included within the desired range of δ phase fraction, namely nomore than 4% and at least 2.5%. It is ascertained that the tensileproperties are distinctly greater than those of products A, B, C and ofthe same level as those of product E. It is also ascertained that thecreep properties of products G, H, M are distinctly better than those ofproducts A, B, C, E whereas the grain size is similar in these products.The fine grain microstructure of products G, H, M allows preservation ofthe fatigue properties obtained with products A, B, C, E, and thesmaller δ phase fraction of products G, H, M allows for improved creepresistance.

The comparison between samples B and P shows that the increase inphosphorus content for a 718 alloy subjected to a reference treatment(a) does not substantially improve creep resistance.

Unexpectedly, the application of a treatment according to the inventionto product M which has a higher phosphorus content (80 ppm), allowed aconsiderable increase in creep lifetimes up to a factor of 4 comparedwith products A, B, C, and also compared with product P whose phosphoruscontent is comparable with that of product M but which was not giventreatment according to the invention.

The combination of added phosphorus with the treatment of the inventiontherefore has a synergic effect which is positive on the creepproperties of the alloy obtained.

The invention targets the maintaining of a residual δ phase fraction(preferably higher than 2.5%) which allows satisfactory ductility to bemaintained at high temperature. If the content of the δ phase is too lowthis has an effect on damage and ductility under tensile testing at hightemperature (650° C. with a strain rate of 10⁻⁵s⁻¹). It is effectivelyascertained that product D whose δ phase content is close to 2% has aductility (elongation at break of 7%) that is much lower than that ofproduct G (elongation at break of 27%) whose δ phase fraction is closeto 3%. This reduced ductility of product D results from intergranulardamage caused by a fraction of δ phase that is too small and distributedheterogeneously.

The influence of the treatments of the invention on the microstructurewill now be detailed.

An examination was made of samples A, B, C, D, E, G, H, I, J, M, O and Pwhich are in 718 alloy and were transformed using thermo-mechanicalschedule N° 2 or n° 3.

FIGS. 4 to 9 are micrographs representing the microstructures:

-   -   of samples A, B, C, D, E, G, H, I, J, M, O and P in their        initial state after thermo-mechanical treatment (FIGS. 4 and 5),    -   of samples D and O after being subjected to heat treatment        comprising only one solution heat treatment step (FIGS. 6 and 7)    -   of samples G, H and M after being subjected to heat treatment        according to the invention (FIGS. 8 and 9).

FIGS. 4 and 5 illustrate the microstructure of samples A, B, C, D, E, G,H, I, J, M, O and P (metallurgical state 1) after being subjected tosub-solvus thermo-mechanical deformation (thermo-mechanical schedule 2or 3). This is a microstructure which shows delta phase Ni₃Nb-δ and/orNi₃Ta-δ at the grain boundaries, but in a manner not uniformlydistributed between the grains.

FIG. 4 shows that the samples have a fine grain of about size ASTM 11,with heterogeneous distribution of the δ phase (black spots at the grainboundaries). After the thermo-mechanical deformation schedule, the δphase percentage is 2.8 to 6% and the grain size is ASTM 10 to 13. Thisgives a very heterogeneous microstructure from these two viewpoints.

FIG. 5 illustrates the microstructure of the samples with greaterenlargement and shows grains whose boundaries are mostly fully devoid ofδ phase (this phase is shown in white in this micrograph).

When a treatment is applied to a sample (sample B) that only comprisesone solution heat treatment step at 970° C. for about 60 minutes, thepercentage of δ phase obtained is 4.7 to 5.5% and the grain size of ASTM11 to 12. The homogeneity of the sample is therefore improved but alarge fraction of δ phase is maintained, for which it is known (seesample B, Tables 1 & 2) that is highly unfavourable for creepresistance.

If the heat treatment applied to a sample (see for example sample O inTable 1) is treatment only comprising one solution heat treatment stepat 1005° C. for about 15 minutes, corresponding to the <<second step>>of the invention, the percentage of δ phase obtained (see FIGS. 6 and 7)is 1.1 to 3.5%, and the grain size is ASTM 5 to 9. The level of δ phaseis therefore reduced, which is in the right direction for creepresistance, but heterogeneous distribution of grain size is observed.This can be accounted for by the heterogeneous grain growth during thisstep resulting from non-homogeneous distribution of the δ phaseinherited from the initial microstructure.

Indeed, and as previously explained, the initial microstructure resultsfrom sub-solvus deformation (state 1), the distribution of the δ phaseis heterogeneous in the initial microstructure. As a result, some grainsmay contain a large quantity of δ phase at the grain boundaries in theinitial structure, whereas other grains only have little or no δ phaseat the grain boundaries (see FIG. 5).

By conducting heat treatment directly at the temperature of the secondstep, without any intermediate temperature hold at the temperature ofthe first step in accordance with the invention, the grains which arenot surrounded with δ phase or which have little δ phase at the grainboundaries will enlarge uncontrollably up to a grain size which mayexceed about ASTM 5-6, whereas the growth of the other grains surroundedby δ phase will be hindered and will give rise to grain sizes close toASTM 9. This heterogeneity of grain size can be clearly seen in themicrographs in FIGS. 6 and 7. The presence, even much localized presenceof ASTM 5-6 grains considerably reduces the fatigue lifetimes.

On the other hand, if some samples (samples G, H and M) are given theheat treatment according to the invention, namely a first step at 980°C. for 60 min and, immediately afterwards, heating as per a ramp of 2°C./min up to a second temperature hold at 1005° C. for 15 min, a δ phaseof 2.9 to 3.5% is obtained with a grain size of ASTM 10 to 12.

The micrographs in FIGS. 8 and 9 show that, compared with the initialstate of the sample:

-   -   the grain size is more homogeneous and remains very fine;    -   the δ phase is now distributed regularly at the grain        boundaries, which efficiently prevents the growth thereof.

By means of the small formation of δ phase precipitates, which leavesthe Nb and Ta elements available in dissolved form, by means of thereduced grain size, of the homogeneous distribution of the δ phase atthe grain boundaries and of the well-adjusted level of the presence ofthis δ phase, creep resistance and tensile strength are improved. It isin particular the fine grain size associated with controlled dissolutionof the δ phase which allows the objectives of the invention to bereached, which are:

-   -   strong fatigue properties, avoiding premature crack initiation        in large grains and giving priority to crack initiation in        niobium carbides;    -   improved yield strength due to more extensive hardening        generated by a larger fraction of hardening phase;    -   a distinct, even considerable, improvement in the creep        resistance of the alloy with sufficient phosphorus content        (sample M).

Once the alloy has been treated according to the invention, thefinishing operations can proceed as is usual in the prior art to obtainthe end part.

In addition, the inventors conducted additional tests on samples ofalloys of type 718Plus and 725, and were able to confirm that theinvention when applied to other nickel-base superalloys having a niobiumand/or tantalum content of more than 2.5%, allowed a marked improvementin their creep resistance and tensile strength.

1. A method for manufacturing a blank of a part in Ni-base superalloycomprising at least 50 Ni in weight percent, comprising: preparing analloy of said superalloy, and conducting heat treatments of said alloy,wherein: said superalloy in weight percentage comprises at least a totalof 2.5% Nb and Ta; heat treatment is applied to said alloy, comprising aplurality of steps distributed as follows: in a first step, during saidalloy is held at between 850 and 1000° C. for at least 20 minutes,precipitating a δ phase at grain boundaries; in a second step, duringsaid alloy is held at a temperature higher than the temperature of thefirst step, allowing partial dissolution of the δ phase obtained at thefirst step, and after the second step to obtain a δ phase quantity ofbetween 2 and 4%, the first and second step being conducted withoutintermediate cooling; and in ageing treatment comprising a third stepand optionally one or more additional steps conducted at a temperaturelower than that of the first step, allowing precipitation of γ′ and/orγ″ hardening phases; the first step being conducted between 900 and1000° C. for at least 30 min, and the second step at between 940 and1020° C. for 5 to 90 min, the difference in temperature between the twosteps being at least 20° C.
 2. The method according to claim 1, whereinan Al content of the alloy is equal to or less than 3%.
 3. The methodaccording to claim 1, wherein a ratio (Nb+Ta+Ti)/Al of the alloy isequal to or more than
 3. 4. The method according to claim 1, wherein agrain size obtained at the end of the alloy treatment is between 7 and13 ASTM.
 5. The method according to claim 1, wherein distribution of theδ phase is homogeneous at the grain boundaries after the ageingtreatment.
 6. The method according to claim 1, wherein after the secondstep a δ phase quantity is obtained of between 2.5 and 3.5%. 7.(canceled)
 8. The process according to claim 1, wherein a changeoverfrom the first step to the second step is performed at a rate of 4°C./min or less.
 9. (canceled)
 10. The process according to claim 1,wherein the alloy comprises by weight: between 50 and 55% nickel,between 17 and 21% chromium, less than 0.08% carbon, less than 0.35%manganese, less than 0.35% silicon, less than 1% cobalt between 2.8 and3.3% molybdenum, at least one of the elements niobium or tantalum, suchthat the sum of niobium and tantalum totals between 4.75 and 5.5% withTa less than 0.2%, between 0.65 and 1.15% titanium, between 0.20 and0.80% aluminium, less than 0.006% boron, less than 0.015% phosphorus,the residual percentage being iron and impurities resulting fromprocessing.
 11. The method according to claim 10 wherein the first stepis conducted at between 920 and 990° C. for at least 30 min and thesecond step is conducted at a temperature of between 960 and 1010° C.for 5 to 45 min.
 12. The method according to claim 11, wherein the totalcontent of Nb and Ta of the alloy is between 5.2 and 5.5%, in that thefirst step is conducted at between 960 and 990° C. for 45 min to 2 h,and in that the second step is conducted at between 990 and 1010° C. for5 to 45 min.
 13. The method according to claim 11, wherein a totalcontent of Nb and Ta of the alloy is between 4.8 and 5.2%, in that thefirst step is conducted at between 920 and 960° C. for 45 min to 2 h,and in that the second step is conducted at between 960 and 990° C. for5 to 45 min.
 14. The method according to claim 1, wherein the alloycomprises a weight content of: between 55 and 61% nickel, between 19 and22.5% chromium, between 7 and 9.5% molybdenum, at least one of theelements niobium or tantalum, such that the sum of niobium and tantalumis between 2.75 and 4% with Ta less than 0.2%, between 1 and 1.7%titanium, less than 0.55% aluminium, less than 0.5% cobalt, less than0.03% carbon, less than 0.35% manganese, less than 0.2% silicon, lessthan 0.006% boron, less than 0.015% phosphorus, less than 0.01% sulphur,the residual percentage being iron and impurities resulting fromprocessing.
 15. The method according to claim 1, wherein the alloycomprises by weight: between 12 and 20% chromium, between 2 and 4%molybdenum, at least one of the elements niobium or tantalum, such thatthe sum of niobium or tantalum is between 5 and 7% with Ta less than0.2%, between 1 and 2% tungsten, between 5 and 10% cobalt, between 0.4and 1.4% titanium, between 0.6 and 2.6% aluminium, between 6 and 14%iron, less than 0.1% carbon, less than 0.015% boron, less than 0.03%phosphorus the residual percentage being nickel and impurities resultingfrom processing.
 16. The method according to claim 1, wherein the alloyhas a weight percent content of phosphorus of more than 0.007%.
 17. Themethod according to claim 1, wherein the first step and the second stepare conducted at sub-solvus temperatures of the δ phase of the alloy,the first step being conducted at a temperature between 50° C. below theδ solvus temperature and 20° C. below the δ solvus temperature, and thesecond step being conducted at a temperature between 20° C. below the δsolvus temperature and the δ solvus temperature.
 18. The methodaccording to claim 1, wherein a temperature of the hot-worked blank partis held constant during at least one of the said steps.
 19. The methodaccording to claim 1, wherein the said third step is conducted atbetween 700 and 750° C. for 4 to 16 h and in that a fourth step isconducted at between 600 and 650° C. for between 4 and 16 h, cooling at50° C./h to +/−10° C./h being carried out between the said third andfourth steps.
 20. The method according to claim 1, wherein between thefirst and second steps, the hot-worked alloy is held at least at oneintermediate temperature between the temperatures of the first andsecond steps for no more than 1 h.
 21. The method according to claim 1,wherein said blank part was prepared in ingot form and then hot-worked.22. The method according to claim 1, wherein said blank part wasprepared using a powder metallurgy method.
 23. A part in nickel-basesuperalloy wherein it is obtained from a blank part manufactured usingthe method according to claim
 1. 24. The part according to claim 23,wherein it is an aeronautic or land-based gas turbine part.